DENSIFICATION AND GRAIN GROWTH

MECHANISMS DURING SPARK PLASMA

SINTERING OF IONIC SOLIDS

A thesis submitted in partial fulfillment

for the award of the degrees

Bachelor and Master of Technology

by

Karthik Akkiraju

MM10B021

Department of Metallurgical and Materials Engineering

Indian Institute of Technology Madras

May 2015

For the love of cup noodles, Coke and coffee.

THESIS CERTIFICATE

This is to certify that the thesis titled ’Densification and Grain Growth Mech-

anisms during Spark Plasma Sintering of Ionic Solids’ submitted by Karthik

Akkiraju (MM10B021), to the Indian Institute of Technology Madras, Chennai for

the award of the degree of B.Tech and M.Tech, is a bona fide record of the research work

done by him under my supervision. The contents of this thesis, in full or in parts, have

not been submitted to any other Institute or University for the award of any degree or

diploma.

Prof. B.S. Murty,

Research Guide,

Dept. of Metallurgical and Materials Engineering

Place: Chennai

IIT-Madras, 600 036

Date: 5th May 2015

Acknowledgments

I shall forever be indebted to Prof. B.S Murty for introducing me to research, having faith

in me for all these years and inspiring me to push the boundaries all along. Dr. Lukas

Bichler, from the University of British Columbia has played an equal part in stitching this

journey together and making Kelowna a place I yearn to be at. This work would have

been incomplete had it not been the engaging discussions over chai with Niraj Chawake,

Dr. Ajit Srivastava, Koundinya NTBN and Dr. Sanjay Kashyap. Thank you very much

for being around on the trips to Ramu and CC. Also, I cannot thank Vishank Kumar,

Audrey Siebert Timmer, Karen Robles and Dr. Apara Ranjan enough for sharing research

woes with Tim Hortons and frequent dinner invites.

I would like to acknowledge the contributions of David Arkinstall (UBCO) and Amit

Sharma (IISc) for helping out with the SEM work. My special appreciation to Anirudha

Karati for helping with the TEM characterization. For their constant advice, Dr. Anand

Kanjarla, Dr. Chinmoy Chattopadhyay, Karthiselva, Pravin from the MME department

at IIT Madras. It was never going to be an individual effort and my research groups at

IITM (Advanced Materials Research) and UBCO have always made me feel at home and

I thank them for bearing me with till I graduate. A huge cheer to my friends Shukla,

Eureka, Mama, Chuddu, Bobo, Belly, Blk, Hotseat, Patil, Chetta Always Wins, and to

you for staying till the end.

Karthik

I

Abstract

In the present work, the densification and grain growth mechanisms during SPS of three

model oxide ceramics: NiO, ZnO and MgO were studied investigating to gain fundamental

understanding of the operative atomisitics. The role of applied pressure, dwell time,

heating rate and electric field was investigated in various temperature regimes.

In the case of NiO nanopowder, the morphological changes and the grain growth mecha-

nism during spark plasma sintering and annealing at different temperatures were studied.

Cuboid shaped grain morphology was observed in the case of pressureless sintering at tem-

peratures above 1000C. The grain growth mechanism involved Brownian motion assisted

grain coalescence via oriented attachment of irregularly shaped nanocrystals, leading to

the formation of nanocubes. The nanocubes subsequently combined to form dense 3-

dimensional microstructure. Similar grain growth mechanism was observed when the

powders were annealed in ambient atmosphere at and above 700 C. Twinning and dis-

location arrays were observed at the resulting interfaces. However, sintering below 1000

C resulted in conventional polyhedral morphology, which was driven by diffusion based

processes. The absence of electric field shifted the density and grain size trajectories to

higher temperatures.

Similarly, SPS of ultra-fine ZnO was carried out in the 600-1000C temperature range

with dwell times of 0-15 min. Between 600-700C, a sintering window was obtained where

grain size and relative density can be controlled. Observation of bridging necks indicated

that localized melting could take place during SPS of ZnO. The presence of liquid film

was related to the role of electric fields via nucleation of Frenkel defects. Formation of

nano grain clusters further indicated the possibility of densification via liquid phase aided

grain rotation. Grain growth and densification mechanisms were evaluated based on the

stress and grain growth exponents. Pore growth and negative shrinkage behavior was

observed at 1000C.

II

MgO nanopowder was sintered in the temperatures range of 900-1400C. Applied pres-

sure was found to lower the sintering temperatures to achieve a given density. A higher

heating rate had a detrimental effect on the densification. Two distinct windows were

observed in the grain size and relative density plots when plotted against temperature.

The sintering behavior was found to be identical in the presence and absence of the elec-

tric field with identical shrinkage curves, grain sizes, relative density and microstructural

features. The grain growth was characterized by the presence of pyramidal-spiral struc-

tures and elongated structures growing from the grain surfaces. It is expected that a

evaporation-condensation mechanism is operative for their growth.

III

Contents

Acknowledgments I

Abstract II

List of Figures VI

List of Tables X

1 Introduction 1

1.1 Sintering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

1.2 Hot Isostatic Pressing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2

1.3 Field Assisted Sintering Techniques . . . . . . . . . . . . . . . . . . . . . 2

1.3.1 Spark Plasma Sintering . . . . . . . . . . . . . . . . . . . . . . . . 2

1.3.2 Flash Sintering . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

2 Literature Survey 5

2.1 Classical Sintering Theory . . . . . . . . . . . . . . . . . . . . . . . . . . 5

Mass Transport Mechanisms . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5

2.1.1 Surface Diffusion . . . . . . . . . . . . . . . . . . . . . . . . . . . 5

2.1.2 Volume Diffusion Lattice Diffusion . . . . . . . . . . . . . . . . . 6

2.1.3 Grain Boundary Diffusion . . . . . . . . . . . . . . . . . . . . . . 6

2.2 Grain Growth . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

2.3 Recent Advances in Sintering Theory . . . . . . . . . . . . . . . . . . . . 7

2.3.1 Effect of Pressure . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

2.4 Material Selection . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9

2.5 Objectives of the Work . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10

2.6 Scope of the Thesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

IV

3 Experimental Details 12

3.1 Raw Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12

3.2 Spark Plasma Sintering and Annealing . . . . . . . . . . . . . . . . . . . 12

3.3 Characterization Techniques . . . . . . . . . . . . . . . . . . . . . . . . . 14

3.3.1 X-ray Diffraction . . . . . . . . . . . . . . . . . . . . . . . . . . . 14

3.3.2 Scanning Electron Microscopy . . . . . . . . . . . . . . . . . . . . 14

3.3.3 Transmission Electron Microscopy . . . . . . . . . . . . . . . . . . 15

3.3.4 Hardness Measurement . . . . . . . . . . . . . . . . . . . . . . . . 15

4 Results and Discussion: Spark Plasma Sintering of nano NiO 16

4.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16

4.2 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

4.2.1 Reduction of NiO . . . . . . . . . . . . . . . . . . . . . . . . . . . 18

4.2.2 Coalescence Driven Grain Growth of Nanocrystals . . . . . . . . . 19

4.2.3 Effect of Boron Nitride layer . . . . . . . . . . . . . . . . . . . . . 23

4.3 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25

4.3.1 Reduction Behavior . . . . . . . . . . . . . . . . . . . . . . . . . . 25

4.3.2 Grain Coalescence . . . . . . . . . . . . . . . . . . . . . . . . . . 25

4.3.3 Effect of Electric Field . . . . . . . . . . . . . . . . . . . . . . . . 30

4.4 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 30

5 Results and Discussion: Spark Plasma Sintering of ZnO 31

5.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31

5.2 Results and Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32

5.2.1 Densification and Grain growth . . . . . . . . . . . . . . . . . . . 33

5.2.2 Observation of sintering necks . . . . . . . . . . . . . . . . . . . . 35

5.2.3 Grain Coalescence . . . . . . . . . . . . . . . . . . . . . . . . . . 38

5.2.4 Effect of Dwell Time: Sintering Analysis . . . . . . . . . . . . . . 39

5.3 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42

6 Results and Discussion: Spark Plasma Sintering of MgO 43

6.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43

6.2 Results and Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43

6.2.1 Decomposition of Mg(OH)2 . . . . . . . . . . . . . . . . . . . . . 43

V

6.2.2 Grain growth and Densification . . . . . . . . . . . . . . . . . . . 45

6.2.3 Microstructural Characteristics . . . . . . . . . . . . . . . . . . . 47

6.2.4 Terraced Oxide Growth . . . . . . . . . . . . . . . . . . . . . . . 50

6.3 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52

7 Conclusions 53

7.1 Recommendations for Future Work . . . . . . . . . . . . . . . . . . . . . 54

References 55

VI

List of Figures

1.1 Schematic of the SPS setup.[3] . . . . . . . . . . . . . . . . . . . . . . . . 3

1.2 (a)Schematic of the flash sintering setup while (b) shows the linear shrink-

age during the flash sintering process.[4] . . . . . . . . . . . . . . . . . . . 4

2.1 Schematic of the possible sintering pathways across the particle interface.[6] 5

3.1 Showing the current path in the case of BN cased NiO (left) and conven-

tional SPS (right). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13

4.1 Dark field TEM image of agglomerate with an average crystallite size of

10 nm while (b) shows the indexed diffraction pattern. . . . . . . . . . . 17

4.2 XRD patterns of the top surface of the sintered pellet showing increased

phase fraction of the reduced Ni layer. . . . . . . . . . . . . . . . . . . . 18

4.3 (a) Line scan across the top surface showing a Nickel layer which gets

oxidized after the thermal etching in (b). The marked line shows the

interface. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18

4.4 (a) Microstructure of the Ni/NiO interface, (b) the growth of the Ni layer

into the NiO particle, (c) reoxidized surface of the Ni/NiO interface, (d)

growth of NiO scales after thermal etching. . . . . . . . . . . . . . . . . . 19

4.5 Fractured surfaces of the samples sintered for 10 min at 50 MPa pressure

(a) at 900C with polyhedral grains and (b) with the faceted cuboidal

morphology developing at 1000C. . . . . . . . . . . . . . . . . . . . . . . 20

4.6 Fractured surfaces of the samples sintered for 10 min at 5 MPa pressure

with (a) cuboidal grains at 1000C, (b) a few larger faceted grains are seen

when sintered at 1050C and (c) cuboidal voids at 1200C with the inset

showing the cubic pores at a higher magnification. . . . . . . . . . . . . . 20

VII

4.7 Plot showing the change in grain size and relative density (inset) as a

function of sintering temperature for the samples sintered with a pressure

of 5 MPa and 50 MPa. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21

4.8 Sequence of the images showing the coalescence mechanism (a) number of

smaller cubes attached to a larger surface, (b) larger cuboidal block and a

cube attachment, and (c) the chain of particles formed as a result of the

coalescence. All images are taken from the powders annealed at 900C for

1hr. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 22

4.9 TEM images of the powders annealed at 700C for 1h with the initial

irregular morphology still being retained, while (b) shows the HRSEM

image evolving cube with some of the retained 111 planes. . . . . . . . 22

4.10 HRTEM of the interface of an impinged particle, (b) showing the indexed

FFT pattern with‘t’ as the twinned spots, and (c) is reconstructed image

of the selected portion showing the twinned region. (d) HRTEM image of

the interface between two particles, (e) reconstructed image using the FFT

showing an array of dislocations and partial dislocations at the interface

with the inset showing a magnified viewed of the defect region. The sample

was annealed at 700 C for 5 min. . . . . . . . . . . . . . . . . . . . . . . 23

4.11 Comparison of the (a) grain size (b) relative density (c) hardness for the

various configuraions studied. . . . . . . . . . . . . . . . . . . . . . . . . 24

4.12 Microstructure of the sintered pellet with BN blocking at (a) 1050C and

(b) 1250C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24

4.13 (a) SEM image of the sample annealed at 1200C showing the initiation

of the attachment process and (b) TEM image of an annealed powder at

700C for 1h showing the cubic building blocks within a large cluster and

a smaller faceted particle is shown to attach to this cluster. . . . . . . . . 28

4.14 Graphical representation of the coalescence mechanism . . . . . . . . . . 30

5.1 SEM micrograph of the raw powder with the inset showing a magnified

image of the nanoparticles. . . . . . . . . . . . . . . . . . . . . . . . . . . 32

5.2 XRD micrographs of the received ZnO powder and pellets sintered at var-

ious temperatures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33

VIII

5.3 (a) Evolution of sintering trajectory and (b) grain size-relative density

trend, as a function of sintering temperature for a dwell time of 5 min. . 33

5.4 Fractured surfaces of ZnO sintered at (a) 600C (b) 800C (c) 1000C. . . 34

5.5 Presence of bridging necks between the particle 600C in (a), (b) and at

1000C in (c). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35

5.6 Line scan across the grain interfaces showing a dip in Oxygen concentration

in (a) while the composition is uniform in (b) for the sample sintered at

1000C for 0 min. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36

5.7 Line scan across the grain interfaces with the necks showing uniform Oxy-

gen concentration in (a) and (b) for the sample sintered at 1000C for 0

min. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36

5.8 Growth of nano grain clusters from (a) 600C to (b) 800C and (c) 1000C. 38

5.9 (a) Change in relative density and (b) grain size as a function of dwell time

at 600C, 700C, 800C and 1000C. . . . . . . . . . . . . . . . . . . . . 39

5.10 (a) Calculation of stress exponent and (b) grain growth exponents. . . . . 40

5.11 (a) SEM images of fractured surfaces of samples sintered at (a) 700C with

15 min dwell and (b) 800C for a dwell time of 10 min. . . . . . . . . . . 41

5.12 Pore enlargement with increase in dwell time at 1000C. . . . . . . . . . 41

6.1 XRD patterns of the heat treated raw powders at various temperatures . 44

6.2 Bright field image of the heat treated powder with the inset showing the

ring pattern. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44

6.3 Shrinkage curves of the samples sintered at 1200C with and without the

application of pressure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45

6.4 Comparative (a) grain size and (b) relative density plot of sintered MgO

with the clouds showing similar trends. . . . . . . . . . . . . . . . . . . . 46

6.5 Fracture surface of the sample sintered at 1200C with a heating rate of

200Cmin−1 and a pressure of 50 MPa. . . . . . . . . . . . . . . . . . . . 47

6.6 Microstrucutural feature of the sintered pellet at 1300C with BN (5 MPa)

showing the neck formation while (b) was sintered at 1300C (50MPa)

showing a stepped surface. . . . . . . . . . . . . . . . . . . . . . . . . . . 48

IX

6.7 Comparison of the microstructure of the sintered samples at 5 MPa with

(a),(b) sintered at 1100C and (c),(d) at 1300C with and without BN

respectively. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48

6.8 Comparison of the microstructure of the sintered samples at 50 MPa with

(a),(b) sintered at 1100C and (c),(d) at 1300C with and without BN

respectively. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 49

6.9 Fractured surface of the sample sintered at 1100C with 5MPa showing

terraced grain structure and (b) showing a pyramid-like terraced structure. 50

6.10 (a) Micron sized rod-like structures while (b)shows the magnified image of

a smaller rod growing from the tip of the terrace structure. The sample

was sintered at 1100C with 5 MPa pressure. . . . . . . . . . . . . . . . . 51

6.11 (a) Triangular islands growing on the surface on the existing micron sized

grains,(b) magnified image of such a structures.The sample was sintered

at 1300C with 50 MPa pressure with BN. . . . . . . . . . . . . . . . . . 52

X

List of Tables

2.1 Comparative properties of the materials used in this study . . . . . . . . 10

3.1 SPS Parameters used in each of the case studies . . . . . . . . . . . . . . 14

XI

Chapter 1

Introduction

Sintering is one of the processing techniques used for ceramic production. In this chapter,

various sintering techniques shall be discussed leading to the debate on spark plasma

sintering and it’s mechanisms.

1.1 Sintering

Sintering is a well established process for producing bulk polycrystalline ceramic mate-

rials. Traditionally sintering has been the thermally driven process to convert the green

bodies into dense ceramic parts (below the melting point). As a result of atomic diffusion,

densification and grain growth occur.

Sintering is classified as solid state and liquid state sintering. Solid state sintering

is driven by the difference in the curvature of the particle surfaces,where the driving

force is given by the Gibbs Thomson equation. Recently, a chemical reaction or external

parameters such as external electric field, magnetic field or pressure are being used to

enhance the sintering kinetics. In addition to this, in liquid phase sintering, the presence

of a small fraction of a liquid phase at the grain boundaries enhances the diffusivities

by 6-8 orders of magnitudes [1]. This help in lowering the sintering temperatures which

can reduce grain growth apart from increasing the economic viability. Apart from the

conventional sintering which uses binders and long sintering times, several techniques

such as Hot Isostaic Pressing (HIP), Field Assisted Sintering Techniques (FAST) such as

Flash sintering and Spark Plasma Sintering, Microwave sintering have been introduced

to reduce the sintering temperatures and have a better control on grain growth and

1

densification. These techniques shall be briefly discussed.

1.2 Hot Isostatic Pressing

In hot isostatic pressing (HIP), a high pressure (100-200 MPa) is applied by means of an

external gas in a sintering furnace upto very high temperatures (2000C). The powder

is encapsulated in a metal container that is placed in vacuum and pressurized. Also,

samples sintered via conventional sintering can be used placed in this furnace for complete

densification. With HIPing near net shape components can be developed, thus eliminating

the need for welding and machining in many cases.

1.3 Field Assisted Sintering Techniques

Apart from applied external pressure, electromagnetic fields can enhance densification

depending on the compatibility of the material with the applied field.

1.3.1 Spark Plasma Sintering

In Spark Plasma Sintering (SPS), sintering is achieved by means of a pulsed electric

field that is applied across the die and the sample along with a pressure. SPS differs

significantly from conventional processes due to the rapid heating rates (>800C min−1)

possible. This is ascribed to the internal Joule heating of the dies and the powders due

to the passage of large currents (thousands of Amperes) rather than the external heating

in conventional techniques [2].

The schematic of the process is shown in Fig. 1.1 [3]. Pressure is applied by means

of stainless steel ramps through which also the current is passed. The temperature

is measured either by a thermocouple (<1000C) that is inserted into the surface of

the die or a pyrometer ( >1000C) focused on the surface of the die. Sintering can

be done either in a vacuum medium or using an inter gas such as argon or helium.

Another feature of the SPS unit is the rapid cooling involved because of the efficient

water circulation. A controlled cooling rate also enables additional heat treatment steps

after the sintering process. Using a computer interface, the temperature and pressure

profiles can be programmed. Linear shrinkage, voltage, current, applied pressure and

2

chamber pressure are constantly monitored during the process. However the scaling of

SPS for industrial applications is an active area of research, with aims to incorporate

larger samples and reduce the power consumption.

Figure 1.1: Schematic of the SPS setup.[3]

Many of the studies in SPS report achievement of near theoretical density at lower

sintering temperatures, shorter times while minimizing grain growth for a wide range of

materials. The underlying mechanisms for such a rapid kinetics shall be discussed in the

next chapter.

1.3.2 Flash Sintering

This recent technique by Cologna and Raj at the University of Colarado involves the

use of extremely large electric fields to sinter ceramics. The schematic of the setup by

Francis et al. [4] is shown in Fig. 1.2(a). A dog-bone shaped green sample is prepared

and hooked using platinum electrodes [4]. The setup is placed in a preheated furnace.

It has been observed that at a particular temperature, remarkable sinterability is seen

in a few seconds. This anomalous densification is invariably accompanied by increase in

the counductivity of the specimen while limiting grain growth. The shrinkage curves for

3

this process are seen in Fig. 1.2(b) where with increasing field strength the flash event is

brought down to lower temperatures. Flash sintering has shown tremendous success with

sintering of nanograined YSZ, Y2O3, Mg doped Alumina and more recently with SiC [5].

The applicability of this process to other non-ionic systems will be an interesting follow

up. Despite this, the low cost of operation has led to the establishment of a company,

Ceram that commercially produces such furnaces.

Figure 1.2: (a)Schematic of the flash sintering setup while (b) shows the linear shrinkage

during the flash sintering process.[4]

4

Chapter 2

Literature Survey

2.1 Classical Sintering Theory

Mass Transport Mechanisms

Sintering occurs due to atomic diffusion at elevated temperatures. To achieve this various

pathways such as surface diffusion, grain boundary diffusion, volume diffusion, lattice

diffusion exist [6]. A schematic for the same is shown in Fig. 2.1.

Figure 2.1: Schematic of the possible sintering pathways across the particle interface.[6]

2.1.1 Surface Diffusion

The surface is source of defects such as extra atoms, vacancies, absorbed species. Due

to this atomic motion occurs between these defects. Atoms diffuse from a convex to a

5

concave curvature to minimize the surface energy. Since this is a re-distribution process,

it is non-densifying.

2.1.2 Volume Diffusion Lattice Diffusion

It is the replacement of atoms with vacancies. Pores, which are large vacancy clusters

are filled with atoms from the surrounding. Such a diffusion from the neck region to

the grain boundary can lead to densification. Pore growth can also occur due to volume

diffusion due to transfer of vacancies across pores.

2.1.3 Grain Boundary Diffusion

This is the most common densifying mechanism for many metals. The activation energy

is lower than that for lattice diffusion but higher than surface diffusion. Due to surface

diffusion, grain coarsening occurs, the driving force for surface diffusion decreases, leading

to the take over by grain boundary diffusion. Here, there is mass transfer from the grain

boundary region to the neck.

Apart from these common mechanisms, plastic deformation, evaporation condensation

have also been reported.

Classical sintering theory divides the densification into three stages : initial, interme-

diate and final. In the initial stage, necks develop at the particle contact points, and the

densfication proceeds via surface diffusion and particle rearrangement. The next stage is

characterized by the change of porosity from the open channels to closed ones. The final

stage, commonly accepted as >92% density involves the elimination of pores at grain

boundaries and grain growth. Mostly grain boundary diffusion and volume diffusion are

active in this regime.

2.2 Grain Growth

Rapid grain growth in nanoparticles is an important consideration for the production of

bulk nanograined materials. The motivation of many sintering techniques is to reduce

sintering temperatures and duration to avoid grain growth. Strategies such as novel

sintering techniques, control of the grain growth mechanism by doping or step sintering

processes are being explored to tackle this problem.

6

The study of grain growth mechanisms can be again divided depending on the size

range. Once the nanosized particles grow into micron sizes, the classical generalized

parabolic equations become applicable which is based on the Ostwald ripening phe-

nomenon in two phase materials. Among various models, the calculation of activation

energies in different regimes and evaluation of the exponent n can give some insights into

the grain growth mechanism. However, given the number of factors that go into such

equations, it often becomes difficult to zero in on a single mechanism.

The driving force for any particle sintering is inversely proportional to the particle

size as given by the equation,

σ = γκ = γ(1

R1

+1

R2

),

where γ is the surface energy of the material, κ is the curvature of the surface, and R1

and R2 are the principle radii of curvature [6]. Hence for nanosized particle, the driving

forces are higher [7]. The initial grain growth mechanisms in the nano range is still

not well understood with surface diffusion, grain boundary migration, coalescence and

reprecipitation.

2.3 Recent Advances in Sintering Theory

With the introduction of new sintering techniques (discussed in the previous chapter)

it has become necessary to develop new theories to explain the anomalous densification

behavior during SPS and flash sintering of nanosized particles. SPS, in particular, has

attracted attention from the sintering community for its rapid heating rate and the con-

tentious role of plasma that is believed to enhance the densification. Flash sintering can

be identified as an extreme case of SPS where very large fields are supplied and the re-

sults from such studies can aid in the understanding of SPS. Similarly, the pressure effects

from HIP or hot pressing studies can be compared to the SPS. Of the several processing

parameters during SPS, the effect of pressure, the rapid heating rate and the electric field

shall be briefly discussed.

2.3.1 Effect of Pressure

With increase in pressure, densification increases invariably. The maximum applied pres-

sure is limited by the strength of the dies. The driving force at a particle interface is

7

given by the equation [8],

µ1 = µ0i ∗ −σnΩI + ezφ,

Where µ0i is the standard reference potential, σn is the normal stress, e is the electron

charge, z is the valence of the diffusing species, φ is the local electric potential and ΩI is

the atomic volume of the diffusing species.

Densification will occur via plastic deformation if the yield strength of the material is

less than the applied pressure. MgO has been densified by SPS via this mode and maps

have been developed for various regimes [9]. Studies have been carried out to use the

models developed for creep and hot pressing for SPS data. Hence, by eliminating the

role of pressure by conducting pressureless sintering, more fundamental insights in the

process can be obtained.

Heating Rate

Rapid heating rate is one of the crucial parameter in SPS. The importance of heating

rate is attributed to the bypassing of the lower temperature surface diffusion regime [10].

The effect of heating rate on densification is also unclear.

Effect of Electric Field

The argument about the presence and absence of plasma has been debated extensively.

While have [11]shown that the plasma was not found even after rigorous experiments.Interestingly,

recently it has been shown by direct visual examination and microstructural character-

ization that a discharge may occur [12]. Chawake et al. [3] have recently shown that in

the case of metal powders such as Fe, Ni and Cu, it is the Joule heating arising from the

contacts of graphite die and punches that is responsible for the densification rather than

the internal current. While other investigations have shown that for conducting powders,

the electric current can enhance defect mobility.

Chen et al. [13] have investigated the role of pulsing during SPS by observing the

growth on the Mo-Si intermetallic layer with varying DC pulsing time. In this indirect

study no effect on the growth rate of the reaction layer was found on changing the pulse

patterns from 2:8 to 8:2 ON/OFF ratio. Although there are several works on the role of

pulsed current during SPS of metals , the fundamental role of the pulsed current during

8

sintering of ceramics in not well established. Shen et al. reported that both densifica-

tion and grain growth were affected by the pulse sequences [14]. At lower temperature

densification was retarded as the ON/OFF ratio increased, at higher temperature grain

size decreased. However, in a recent study no effect on the pulse pattern during sintering

of Al2O3 was found [15]. Pulse sequence ranging from 2:2 to 48:2 had no effect on the

densification of spark plasma sintered yttria stabilized zirconia [16].

The case for non-conducting materials or ionic ceramics is inconclusive with several

conflicting results on the role of grain boundary mobility. While very large electric field

have shown to enhance densification, the role of weak fields in SPS is currently being

debated. Ghosh has shown that weak DC electric fields can retard grain growth which

is contrary to the normal expectation [17].While YSZ has been used as model ceramic

for most such experiments, the role of weak fields in other ceramics is not understood.

Narayan [18] [19] strengthened the cause for electric field effects by proposing a surface

melting phenomenon and a new grain growth model based on his observations on MgO.

Although the electric field and current cannot be explicitly controlled in SPS, their

role in the sintering processes needs to be revisited.

2.4 Material Selection

By understanding the role of the applied pressure and the electric field during spark

plasma sintering of oxides in enhancing the densification during SPS, a better under-

standing of the atomistics during SPS can be gained.

During their in-situ experiments of field assisted sintering of Ni nanoparticles which

had a thin NiO layer, Bonifacio et al. observed the time dependent breakdown of the NiO

layer to given an Ni layer. Metallic Ni necks were formed between the particles, which

was followed by surges in the current value [20]. This experiment using an STM tip in

a TEM machine gave a fundamental insights of the field effects during ECAS of NiO.

Interestingly, Chaim and Bar-Hama [21] densified nano NiO and associated the rapid

densification to the creep and grain boundary sliding. Hence, these two studies provided

information about the possible mechanisms for NiO, both of which are possible during

SPS.

9

Similarly, Narayan developed his ’Surface Melting’ [19] [18] theory based on his obser-

vation during field treatment of MgO single crystals. In such two electrode experiments,

the MgO single crystals were found to fail by a Joule heating along dislocations and va-

cancy clusters [22]. It has been proposed that this heating can cause rapid densification

in FAST. In addition to this, sintering maps developed for SPS of MgO attribute the

rapid densification to the plastic deformation [9]. Here again, the proven role of electric

field in MgO, provides an opportunity for investigation into the operative mechanisms

during SPS of MgO.

Langer et al. have concluded that there is no difference in the densification mecha-

nisms and grain growth behavior between ZnO processed by HP and SPS, suggesting that

the current plays no significant role during the sintering process [23]. However, Misawa

et al. clearly observed particle boundaries in the case of DC pulse-free SPS processing,

while the boundaries were absent when DC-pulses were active [24]. Further, during SPS

of nanocrystalline ZnO at 500 C, an amorphous intergranular ZnO layer was observed,

presence of which resulted in abnormal electrical conductivity of the as-sintered mate-

rial [25]. These contradictory results indicate that there is still not a clear understanding

on the role of electric field during sintering of ZnO.

A comparison of the properties of three oxides is shown in Table. 2.1.

Table 2.1: Comparative properties of the materials used in this study

Material Melt. Temp.(C) Dielectric Constant Crystal Structure Band Gap(eV)

NiO 1955 9-12 Rocksalt 3.4

ZnO 1975 10 Wurztite 3.3

MgO 2852 10 Rocksalt 7.8

2.5 Objectives of the Work

The objectives of the present work are:

• To densify three model oxide systems: NiO, ZnO and MgO nanopowders using

Spark plasma sintering.

• Study of the role of process parameters by systematic deconstruction of each factor

by means of pressureless sintering and electric field blocking experiments using

10

Boron Nitride (BN).

• To evaluate the grain growth and densification mechanisms using microscopic tech-

niques and analytic models in each of such cases shall be probed.

• To identify the processing parameters among the pressure, heating rate, electric

field and dwell time that are significant for the densification of ceramic oxides.

2.6 Scope of the Thesis

Limited experimentation of the field effects during sintering of NiO, MgO and ZnO,

and the availability of classical densification mechanisms during SPS, provides a case

for the investigation of these oxide materials with emphasis on the role of pressure and

electric fields. This approach has been attempted by conducting conventional pressure

assisted sintering along with pressure less experiments and field blocking experiment

using hexagonal BN. The Results and Discussion of the thesis are organized separately

for the oxide systems with an underlying theme to gain a common understanding of the

process. With similar experimental design across the three materials, the conclusions

chapter compares the results and presents a common understanding.

11

Chapter 3

Experimental Details

3.1 Raw Materials

Raw powders of NiO (Inframat Advanced Materials, USA) , ZnO (Fischer Scientific,

USA), MgO (Inframat Advanced Materials, USA) were purchased and used directly for

the experiments except in the case of MgO where prior heat treatment was necessary to

eliminate the formed Magnesium Hydroxide.

3.2 Spark Plasma Sintering and Annealing

Different SPS parameters had to be employed for different materials, owing to the differ-

ences in the melting point. ZnO and NiO were sintered in a 20 mm diameter graphite

die lined with grafoil using the Model 10-3 (10 ton/3000 A) Thermal Technology SPS

unit while MgO was sintered using the Dr. Sinter SPS-5000 machine. For all sintering

experiments, a 5 MPa pre-load was applied to achieve a green pellet density of 45%.

Pressureless sintering was carried out with the 5 MPa load maintained throughout the

entire SPS run.

For the field blocking experiments, the ceramic samples were sandwiched between lay-

ers of hexagonal Boron Nitride (BN) as shown in Fig. 3.1. For pressure assisted sintering

experiments, 50 MPa load was used for all the three oxides. The optimized temperature

and dwell times investigated are reported in the table below. For temperatures inves-

tigated below 1000C, thermocouple inserted into the bottom die was used to monitor

the temperature. For temperatures above 1000C, a pyrometer focused on the surface of

12

the die was used. In this case, temperatures below 600C cannot be measured and the

heating rate in this range is uncontrollable. Real time sample shrinkage information was

obtained from the sample dimensions and the SPS ram displacement data collected by

the data acquisition unit.

Figure 3.1: Showing the current path in the case of BN cased NiO (left) and conventional

SPS (right).

The density of the sintered pellet was measured using the Archimedes principle.

ρ =ma

ma −mw

∗ ρH2O

Where ma and mw are the weight of the sample in air and in de-ionized water respectively.

The instantaneous relative density of the pellet was estimated using the relation,

ρi = ρf ∗hfhi

Where ρi is the instantaneous density, ρf is the final density and hf , hi are the corre-

sponding heights of the specimens.

The densification strain rate behavior was studied via Eq. 1,

1

ρ

dρ

dt= HD

φPa

RTGn

The equation governing grain growth can be written as Eq. 3, simplified from the

pore drag model by Kang [1],

Gn −Gn0 = kt

13

Where G0 is the average grain size at time t=0, k is the rate constant and n is the grain

size exponent. If G0 is small when compared to G, the slope of the lnG vs lnt plot yields

the grain size exponent n.

Table 3.1: SPS Parameters used in each of the case studies

Material Temperature( C) Dwell (min) Heating Rate (C/min)

NiO 900-1300 10 200

ZnO 600-1000 0-15 200

MgO 900-1400 3 200,100

In order to compare the grain growth occurring during sintering process and annealing

process, annealing experiments were conducted by placing the raw powders in a muffle

heat treatment furnace by directly loading them at the set temperature.

3.3 Characterization Techniques

3.3.1 X-ray Diffraction

X-ray diffraction (XRD) experiments were conducted using the X’Pert Pro (PANalytical)

X-ray Diffractometer in Bragg-Brentano geometry equipped with X’Celerator detector

with Cu-Kα radiation (45 kV, 30 mA) in the 30-90 2θ range.

3.3.2 Scanning Electron Microscopy

Powders (i.e. as-received and annealed nanopowders) along with fractured and polished

surfaces of the as-sintered ceramic samples were observed using a Tescan MIRA3 XMU

equipped with and FEI QUANTA 400 Field Emission Gun equipped scanning electron

microscopes (SEM). For grain size measurements, the sintered samples were cut, polished

to 1 µm and then thermally etched at 100C below the sintering temperature to reveal

the grain boundaries and the largest axis of the grain polygon was measured.

14

3.3.3 Transmission Electron Microscopy

Transmission electron microscopy (TEM) was carried out using a Tecnai T-20 unit op-

erating at 200 kV, while high resolution TEM (HR-TEM) imaging was carried out with

Tecnai F-30 field Emission TEM operating at 300 kV to observe grain interfaces in as-

sintered samples. Fast Fourier Transform (FFT) and reconstructed lattice images using

the Digital MicrographTM software (Gatan Microscopy Suite) were used to analyze the

HRTEM images. For the powders, the samples were deposited on a copper grid after dis-

persion in acetone. For the sintered samples, ion milling was carried out after dimpling.

Additional cross-sectional method was usd to prepared TEM samples as the ceramic

samples were too brittle for conventional techniques.

3.3.4 Hardness Measurement

Hardness was evaluated using the Vickers micro hardness indenter with a load of 1N for a

dwell time of 15 sec. An average of 20 measurements was used to estimate the hardness.

15

Chapter 4

Results and Discussion: Spark

Plasma Sintering of nano NiO

4.1 Introduction

Processing of nanocrystalline materials is fundamentally different from sintering of con-

ventional materials, since the larger surface area of nanocrystalline powders enhances

the sintering kinetics [7]. Ostwald ripening (OR) and oriented attachment (OA) are

the mechanisms that are usually suggested to explain the sintering and grain growth in

nanocrystalline materials [26]. OR is a classical grain growth mechanism where larger

grains grow at the expense of smaller ones, and the driving force for grain growth comes

from the excess surface energy, with the growth rate given by the Lifshitz-Slyozov-Wagner

(LSW) equation [27]. In contrast, OA grain growth model proposed by Penn et al. [28]

for nanocrystalline systems suggests that particles rotate until an epitaxial or twin con-

figuration is reached. Though both grain growth mechanisms involve reduction in surface

energy, the ability of grains to rotate is inversely proportional to their grain size [29]. As

a result, the OA mechanism is expected to be prominent in nanocrystalline materials.

The ability of the SPS process to rapidly sinter nanopowders with minimal grain

coarsening can create new pathways for control of grain growth kinetics and grain mor-

phologies [30]. Owing to the high surface energy of the nanograins, the anisotropy of the

surface energy has to be considered during sintering. Few studies in the literature have

reported the impact of surface energies on the grain growth behavior of nanomaterials

during sintering. A number of published investigations attempting to describe the grain

16

growth mechanisms in nanograined materials have been carried out with colloidal systems

during hydrothermal growth. In such systems, along with the temperature, surfactants

which adsorb differently on planes with different atomic density may affect the relative

growth rate of various planes [31]. However, in solid state sintering processes such as

SPS, rapid heating rate can also be a useful parameter to control the grain morphology

and the growth mechanism. Hu et al. [30] have reported grain growth by a novel multi or-

dered coalescence mode of cubic nanocrystals of SrTiO3, where microstructure evolution

was governed by the rate of heating.

In the present work, nanocrystalline NiO was used as a model system to study the

grain growth behavior during SPS of nanomaterials. NiO is an important material with

a potential for thermoelectric material applications [32], anodes in Lithium-ion batteries

[33] and supercapacitors [34], where the interplay of porosity and grain size can be used

to tune the ceramic’s functional properties.

4.2 Results

TEM micrographs revealed that the as-received NiO powders formed nanocrystalline

agglomerates. The average crystallite size was estimated to be 10 nm from the dark

field TEM images(Fig. 4.1(a)), which correlated well with the crystallite size of 12 nm

obtained from XRD analysis (Fig. 4.2). From Fig. 4.1(b), the ring pattern was the

identified as a rock salt structure of NiO.

Figure 4.1: Dark field TEM image of agglomerate with an average crystallite size of 10

nm while (b) shows the indexed diffraction pattern.

17

Figure 4.2: XRD patterns of the top surface of the sintered pellet showing increased

phase fraction of the reduced Ni layer.

4.2.1 Reduction of NiO

The sintered samples had a thin metallic layer on the outer surface. Elemental analysis

of this layer showed that the layer consisted of metallic elemental layer. After thermal

etching the nickel oxidized to form Nickel Oxide againas (Fig. 4.3(a,b)). The difference

in the composition in both these cases is illustrated in the line scans in Fig. 4.3. The

oxidized micro-structure consisted of beautiful scales of NiO, as seen in Fig. 4.4

Figure 4.3: (a) Line scan across the top surface showing a Nickel layer which gets oxidized

after the thermal etching in (b). The marked line shows the interface.

18

Figure 4.4: (a) Microstructure of the Ni/NiO interface, (b) the growth of the Ni layer

into the NiO particle, (c) reoxidized surface of the Ni/NiO interface, (d) growth of NiO

scales after thermal etching.

4.2.2 Coalescence Driven Grain Growth of Nanocrystals

In the case of pressure assisted sintering, polyhedral grains with average size of 0.386 nm

and 0.473 nm were observed when the NiO was sintered at 900 and 1000C, respectively.

Pores were located either at the triple junctions or at the grain boundaries in accordance

to classical sintering theory (Fig. 4.5). As the sintering temperature increased to 1100C,

the polyhedral grains and spherical pores were no longer observed. Instead cuboidal pores

were seen (Fig. 4.5(b)). Fractured surface of the pressureless sintered samples at 1000C

showed cubic morphology of the grains with an average grain size of 68 nm and relative

density of 79% as seen in (Fig. 4.6(a)). At 1025C, the relative density increased to

88%, while the average grain size increased to 100 nm. At 1050C, the density increased

to 90%, and the microstructure was marked with the presence of larger micron sized

cuboidal grains along with nanosized cubes in their vicinity (Fig. 4.6(b)).

19

Figure 4.5: Fractured surfaces of the samples sintered for 10 min at 50 MPa pressure (a)

at 900C with polyhedral grains and (b) with the faceted cuboidal morphology developing

at 1000C.

Figure 4.6: Fractured surfaces of the samples sintered for 10 min at 5 MPa pressure with

(a) cuboidal grains at 1000C, (b) a few larger faceted grains are seen when sintered at

1050C and (c) cuboidal voids at 1200C with the inset showing the cubic pores at a

higher magnification.

A further increase in the sintering temperature showed that the average grain size

remained unaffected until 1075C, at which point an abrupt increase in grain size was

observed (from 110 to 1400 nm) at temperatures between 1075 and 1100C) as shown in

Fig. 4.7. After 1100C, the grain size did not increase rapidly and reached 2.25 µm at

1200C. Transgranular fracture of the sample sintered at 1200C revealed cuboidal pores,

as observed in Fig. 4.6(c).

20

Figure 4.7: Plot showing the change in grain size and relative density (inset) as a function

of sintering temperature for the samples sintered with a pressure of 5 MPa and 50 MPa.

The observation of the cuboidal morphology of the grains and the pores in the pres-

sureless sintering experiments at and above 1000C and the change in the pore mor-

phology from spherical to cuboidal between 1000 and 1100C for the pressure assisted

case indicates the presence of an activation barrier for this morphological transformation.

This transformation appears to be independent of the grain size or the relative density

as the pressure assisted samples had much higher densities and grain sizes when com-

pared to the pressure-less case (Fig. 4.7). The presence of polyhedral shape at lower

temperature and the formation of nanocubes at higher temperature could be the result

of different grain growth mechanisms active at different temperature ranges. In order to

verify whether the grain growth mechanism is a feature of any spark/plasma event of the

SPS unit or an inherent nature of the nanocrystalline system, the microstructure of the

annealed as-received powders was examined.

To further investigate the facet development and coalescence behavior, the powders

were observed in a TEM. The powders annealed at 900C showed nanocubes with a

size range between 3-200 nm. The attachment of the smaller nanocubes to the existing

grains provides evidence that the stable larger grains can act as base for subsequent

growth (Fig. 4.8(a)). At some locations, a few cuboids and nanopores were observed

in the TEM images, indicating that the nanocubes were formed through assembling of a

number of nano grains. This could lead to the formation of a continuous chain of particles

(Fig. 4.6(a) and Fig. 4.8(b,c)).

21

Figure 4.8: Sequence of the images showing the coalescence mechanism (a) number of

smaller cubes attached to a larger surface, (b) larger cuboidal block and a cube attach-

ment, and (c) the chain of particles formed as a result of the coalescence. All images are

taken from the powders annealed at 900C for 1hr.

Akin to the sintered samples, the annealed powders also had the nanocube grains and

the microstructure appeared to consist of nanocube shaped building blocks. The cubic

grains first appeared after annealing at 700C for 1hr, but a few nanocrystals remained

untransformed, as seen in Fig. 4.9(a). Temperature driven nanograin coalescence appears

to be an inherent character of the nanocrystals. Ultra high resolution SEM imaging of

the cube structures revealed that in some cases the edges had triangular facets, while in

some locations edges were perfectly formed, as observed in Fig. 4.9(b).

Figure 4.9: TEM images of the powders annealed at 700C for 1h with the initial irregular

morphology still being retained, while (b) shows the HRSEM image evolving cube with

some of the retained 111 planes.

To investigate the interfaces formed during the attachment process, HRTEM imaging

was carried out on both the annealed powder and as-sintered NiO. FFT and reconstruc-

tion of the square region marked in Fig. 4.10 of the pellet sintered at 1075C showed

lattice fringes with an interplanar spacing of 0.21 nm. Indexing using standard database

22

(JCPDS PDF: 47-1049) identified these interfaces as the 200 family of planes of NiO.

Hence, the nanocubes can be assumed to be made up of 100 planes.

Figure 4.10: HRTEM of the interface of an impinged particle, (b) showing the indexed

FFT pattern with‘t’ as the twinned spots, and (c) is reconstructed image of the selected

portion showing the twinned region. (d) HRTEM image of the interface between two

particles, (e) reconstructed image using the FFT showing an array of dislocations and

partial dislocations at the interface with the inset showing a magnified viewed of the

defect region. The sample was annealed at 700 C for 5 min.

The coalescence of the 100 planes has been investigated by intentionally arresting

the annealing process and then examining the microstructure. For powders annealed at

700C for 5 min and the sample sintered at 1075C for 10 min, twin configuration was

observed at some locations (Fig. 4.10(b)). Extra spots in FFT pattern (inset of Fig.

4.10(c)) confirmed the presence of the twin. Coalescence along 111 contact plane led

to twin formation in the both the annealed and the sintered specimen with [011] twin

axis, as indexed in Fig. 4.10(b). At other interfaces, reconstructed images using the FFT

patterns revealed an array of dislocations along the interface of two coalescing particles

(Fig. 4.10(d,e)) while at other locations the interfaces were defect free.

4.2.3 Effect of Boron Nitride layer

A clear distinction in the relative density, grain size and hardness was observed in the

case of electric field comparison studies. For a given temperature, the BN samples always

23

had a lower grain size, relative density and hardness, as inferred from Fig. 4.11.

Figure 4.11: Comparison of the (a) grain size (b) relative density (c) hardness for the

various configuraions studied.

The microstructural difference between the samples sintered with and without the BN

casing as shown in Fig. 4.12(a,b). In the case of the BN samples, the microstructure at

1050C resembles that of the conventionally sintered samples, with no trace of the cubic

morphology. The grains begin to develop polyhedral facets instead of the 100 seen in

the pressureless case. As the sintering temperature increased to 1250C, a few cuboidal

pores were observed along with the spherical ones, while the grain morphology remained

polyhedral.

Figure 4.12: Microstructure of the sintered pellet with BN blocking at (a) 1050C and

(b) 1250C.

24

4.3 Discussion

4.3.1 Reduction Behavior

The presence of the Ni rich layer indicates that NiO has undergone reduction. There are

several reasons for this,

• In-situ reduction due to the electric field. Holland et al. [20]have reported that Ni

can form during in-situ TEM experiments when electric field is applied via an STM

tip.

• Reduction due to reaction with graphite.

• Decomposition to Ni due to presence of vacuum in the chamber, following the

Ellingham’s diagram.

As the Nickel layer was present in the case of the BN-NiO case as well, the loss of oxygen

can be ascribed to the low oxygen pressure in the chamber. However, the contribution

from other factors have still to be evaluated.

4.3.2 Grain Coalescence

The results suggest that the microstructure of the SPS and annealed samples is similar.

Thus, it is likely that formation of grain interfaces and hence the grain growth mecha-

nism followed a similar path in both the cases. It has to be noted that grain growth in

nanograins cannot be characterized by a single growth mechanism, since the attachment

and coarsening of the nanograins may be occurring simultaneously. Extensive in-situ

TEM experiments are necessary to unambiguously determine the grain growth mecha-

nism [35]. In this work, OA was assumed to be the dominant mechanism based on the

microstructural observations.

If the nanocrystals are treated as individual entities, understanding of their growth

mechanism must consider the driving force for an oriented attachment growth, the shape

of the final assembly and the mechanism of attachment [36]. The mode of growth by

crystal attachment has been well established during hydrothermal growth of nanoparticles

in colloidal systems and has been termed as grain rotation induced by grain coalescence

(GRIGC) [37].

25

Gong et al. have suggested that OA occurs in a stepwise process where the initial

nanocrystal aggregates rotate to facilitate the collision of high energy surfaces followed

by the removal of surfactants and formation of coherent interfaces [38]. By adjusting the

reaction temperature and annealing time of FeS2 nanoparticles they have obtained cubic

and sheet morphologies. In the present work, particle clustering in the as-received powder

and the annealed powders, along with the appearance of cuboidal grains at 1000C in

the sintered specimens, suggests that a critical temperature needs to be achieved before

sufficient kinetic energy is available for the nanocrystals to overcome attractive forces

within the raw powder agglomerate. For this to occur, sufficient Brownian motion of the

crystals has to be ensured, especially for solid state growth where surfactants or a liquid

medium which can aid attachment are not available. At high temperatures (above a

critical temperature, Tc), the rapidly moving nanoparticles may collide with a favorable

crystallographic plane and then coalesce [38]. The observed initial particle clustering

(where common crystallographic orientation may not be achieved) could aid in the OA

process as the cluster in Fig. 1 may act as the precursor for subsequent formation the

cubic grains [16]. The formed clusters then transform into larger single crystals, as seen

in the sintered sample in Fig. 4.5 and in the sample annealed at 900C (Fig. 4.9). The

formation of single crystal has been confirmed from the electron diffraction patterns.

The starting nanocrystals form clusters and assume a shape depending on the surface

energies of the planes. The equilibrium shape is given by Wulff plot [39], where the higher

surface energy faces grow faster and finally are not present in the equilibrium morphology.

Lee et al. [31] have proposed that intrinsic surface energy of crystallographic faces, amount

of adsorbing species, sufficient time for growth of thermodynamically stable structures

and molecular precursors are the four factors that govern the growth behavior in the case

of hydrothermal growth. In the case of solid state sintering without any sintering aids,

the surface energies of the crystallographic faces of the starting particles and sufficient

time at the growth temperatures are the applicable factors. Hence, the competition

between these thermodynamic and kinetic factors has to be studied to understand the

grain growth behavior.

The temperature dependence of the growth of various planes has to be considered,

taking into account that NiO crystallizes in the rock salt configuration with O2− at the

(0,0,0) position and Ni2+ at the (12

,12

,12) positions in the investigated temperature range.

26

Typically for grain growth of rock-salt crystals, a polyhedron is assumed to be formed

initially at low temperature and is bound by six 100 and eight 111 faces. The final

crystal morphology is dependent on the rate of evolution of these faces [40]. Theoretical

calculations have shown that 100 facet in NiO has the lowest surface energy (0.958

J m−2) while the 110 is the second lowest (1.26 J m−2) [41]. HRTEM images of the

NiO cubes showed that the interplanar spacing was 0.21 nm, corresponding to the 200

planes. Hence the surface of the NiO cube was formed by six 100 faces. If sufficient

thermal energy (kT) is provided, rapid growth occurs in the <111 >direction and cubic

morphology evolves. In case sufficient energy or time are not available, truncated octa-

hedron can result due to the slower growth kinetics of the 111 planes when compared

to the 100 faces. While for the pressureless sintered NiO specimens there was sufficient

thermal energy for the 111 planes to grow, in the pressure-assisted case, the polyhedra

remained. Hence the triangular planes observed in Fig. 5(b) are the 111 planes that

have not yet fully grown out.

Chaim and Bar-Hama [21] have previously densified nano-NiO using SPS with the

powders from the same manufacturer that were used in this study. Interestingly, they

have not reported the formation of any nanocubes and their TEM micrographs reveal only

conventional polyhedral structures at 900C. In the present study, when a pressure of 50

MPa was applied at 900C, NiO grains with polyhedral morphology were obtained in Fig.

4.5. Sahoo et al. also sintered nanocrystalline NiO using SPS in the temperature range of

400-800C and have reported clean grain boundaries [42]. Diffusion-driven mechanism can

be responsible for such grain growth at lower temperatures. However, there was a shift

in the temperature associated with the evolution of nanocubes, with the highest being

for pressure–assisted, then pressure-less sintered and finally for annealed samples. This

can be attributed to the steric hindrance (function of relative density) for the nanocrystal

movement [43].

The subsequent microstructure evolution is dependent on the attachment of the

nanocubes. For particle coalescence, attachment may begin at the edges and corners

(as seen in Fig. 4.9(a)), where the electric double layer that prevents coalescence is

absent [44]. The initial attachment is followed by the formation of the quasi sintering

neck, as seen in Fig. 4.13(b)). Formation of such necks was observed in the case of tin

oxide [37]. The kinetics of subsequent coalescence were described with the A1+A1 model,

27

where starting nanocrystals (A1 blocks) combined with each other by sharing a face and

decreasing the surface energy during coalescence. The blocks formed in Fig. 4.13(b)

appear to be the result of such a process. However, as seen in Fig. 4.9(c), a continuous

coalescence of different sizes of crystals was also observed. Therefore, the overall growth

process is likely described by the Ai+Ai model, where the attachment of primary crystals

to an existing higher-order crystal substrate takes place. Hence, the abrupt jump in the

grain size at 1100C may be associated with the primary cubes coming together to form

larger crystals. Similarly variation in grain size was observed in the case of SrTiO3 [30],

CeO2 [45] and it was suggested that the larger grains were formed by aggregation with

smaller ones. Interestingly, the assemblage of the smaller nanocubes could not be cap-

tured in the present work likely due to the narrow window of the exaggerated growth.

In-situ TEM studies by Theissmann et al. [43] revealed that the coalescence process oc-

curs in less than 1/25 s. Thus, the initial grain clustering and the sudden surge in grain

size indicate that microstructure evolution progressed via a continuous buildup of the

nanocubes.

Figure 4.13: (a) SEM image of the sample annealed at 1200C showing the initiation

of the attachment process and (b) TEM image of an annealed powder at 700C for 1h

showing the cubic building blocks within a large cluster and a smaller faceted particle is

shown to attach to this cluster.

Rapid heating in the SPS enables the initial nanocrystals to reach higher temperatures

without any coarsening or necking in order to promote the formation of nanocubes. Hu

et al. observed that multi ordered coalescence of cubic nanocrystals was promoted by

a high heating rate. In contrast, at lower heating rate the ordered coalescence was

significantly reduced due to the formation of necks at lower homologous temperatures.

28

These studies suggest that a critical temperature should be reached rapidly in order

to facilitate grain motion. In this study, a heating rate of 200C min−1 was applied

during SPS processing, which could have enabled sufficient Brownian motion for random

collision and coalescence of the nanoparticles. The pores observed in Fig. 4.5 were likely

the result of improper building of the microstructure due to excessive kinetic energy of

the nano sized units. Similarly, annealing of the nanopowders at a high temperature

enabled the activation of such grain motions and morphological transformations, further

establishing the significance of the heating rate in SPS. Such a pathway is active only when

sufficient Brownian motion is available for the nanocrystal motion without significant

grain coarsening.

Rotational activity continues until a common low energy configuration is obtained,

which is governed by either a parallel epitaxial relationship or a twin configuration (Fig.

8(b)). Twinning was observed in the case of other oxide particles, such as CeO2 and

TiO2 [46] during coalescence of 111 and 112 planes, respectively. In the present work,

attachment involved the contact plane 111. However, incomplete attachment process

can lead to the formation of line/planar defects which are local minima in the energy

configuration [47]. The presence of large number of line defects, planar defects and pores

cannot be explained by OR. In the model developed by Huang et al., it has been shown

that during coalescence dislocations arise as a result of the large strain mismatch and the

critical grain size for dislocation generation is inversely proportional to the strain [48].

Dislocations have been found during SPS of Al2O3, which were attributed to the strain

created at the grain boundaries during the sintering process. However, these dislocations

disappeared after annealing [49]. In the present work, the presences of such defects in the

annealed powders as well as the sintered samples indicate that the coalescence process

might be a source of these defects.

Bulk material with tailored interfacial defects, morphologies and porosity are impor-

tant for energy, sensing and catalytic applications. Along with the conventional use of

SPS for densifying ceramics, the results of this study indicate that it can also be used as

important tool to control the morphology as well the interface defects when the pressure

and temperature regimes are carefully chosen.

29

4.3.3 Effect of Electric Field

The role of the BN layer can be attributed to two factors, namely as an electric field barrier

and a thermal barrier. While previous studies and our own results on MgO have shown

that BN is a reliable material for such experiments without inducing large temperature

gradients, the cause for the shift in the morphologies and the data obtained in Fig. 4.11

and Fig. 4.12 have to be attributed to the field itself. Tang et al. [50] have shown that

there is difference in the hardness when a electric field was applied. The lowering of grain

size when a field electric field was applied has been investigated previously. Although,

we have not completely ruled out the presence of the temperature gradient, the results

here are an indication that the electric fields may enhance the densification in SPS.

4.4 Conclusions

A novel cubic growth mechanism of nanograins was reported for the first time during

sintering of NiO. For the 3-D microstructure construction, cubic nanostructures act as

building blocks. As evident from TEM analysis, larger cubes, due to their relative stability

can act as a substrate to which smaller cubes attach. A schematic of the process is

illustrated in Fig. 4.14. Crystal growth via the OA mechanism resulted in the formation

of building blocks, along with defects at the interfaces. The observation of a similar

growth behavior in the annealed powders and SPS samples suggests that the coalescence

behavior may not be a feature of SPS, but a characteristic grain growth mechanism of

nano sized particles at high temperature. The electric field has been shown to influence

the morphology and the properties, however the factor of temperature gradients needs to

verified.

Figure 4.14: Graphical representation of the coalescence mechanism

30

Chapter 5

Results and Discussion: Spark

Plasma Sintering of ZnO

5.1 Introduction

Zinc oxide (ZnO) is a wide band semiconductor with applications in the electronic, optical

and bio medical industry. Grain growth and densification studies of pure and doped ZnO

processed by conventional sintering, hot pressing (HP)and SPS, have been previously

reported [23]. A fundamental study by Holland et al. to evaluate the thermal contri-

butions due to rapid heating rates and athermal contributions arising from the applied

field during sintering of ZnO led to the conclusion that electric fields coupled with the

high heating rate enabled enhanced sinterability of ZnO [51]. However, Langet et al. [23]

showed that electric field did not influence the sinterability of ZnO. Such contradictory

conclusions on the role of electric field during sintering of ZnO in literature could be due

to the differences in the range of the strength of the applied electric field, particle size of

the starting ZnO powders and investigated temperature ranges.

The synergetic effect of electric fields, rapid heating rates and applied pressure make

SPS a difficult process to analyze from the existing framework and much uncertainty

remains on the operative mechanisms [52]. Therefore it is of interest to reinvestigate the

microstructure of ZnO ceramics processed via SPS, with an aim to gather evidence of

athermal contributions (if any), in order to yield a better understanding of the atomistic

phenomena during consolidation via SPS. As traditional sintering theories have been

insufficient to explain the rapid sintering associated with SPS, alternate mechanisms

31

should be explored along with careful evaluation of the microstructure sintering analysis

to understand the operative mechanism. In the current study, densification behavior and

microstructure evolution of ultra-fine ZnO powder processed via SPS was studied and

the grain growth mechanisms were evaluated critically with consideration of liquid-like

neck formation and nano grain rotation.

5.2 Results and Discussion

The as-received ZnO powder had a bimodal particle size distribution with faceted particles

in the range of 150-200 nm, along with fine spherical particles of 20-30 nm in diameter

(Fig. 5.1).

Figure 5.1: SEM micrograph of the raw powder with the inset showing a magnified image

of the nanoparticles.

XRD analysis confirmed that the starting powder and all the sintered samples had

peaks corresponding to hexagonal Wurtzite ZnO phase (Fig. 5.2).

32

Figure 5.2: XRD micrographs of the received ZnO powder and pellets sintered at various

temperatures.

5.2.1 Densification and Grain growth

Effect of Temperature

As the sintering pressure increased to 50 MPa from the pre load of 5 MPa, the rela-

tive density increased from 45% to 62% and remained constant until shrinkage began at

350C. The development of the relative density in Fig. 5.3(a) with respect to tempera-

ture indicates that the densification rate reached a maximum in the 600-700C sintering

temperature range.

Figure 5.3: (a) Evolution of sintering trajectory and (b) grain size-relative density trend,

as a function of sintering temperature for a dwell time of 5 min.

33

The difference in the kinetics of densification and grain growth were reflected in the

nature of porosity and grain morphology. At 600C, the microstructure contained open

interconnected porosity, while the grain surfaces were curved, with an average grain size

of 290 nm (Fig. 5.4(a)). With further increase in the sintering temperature, a gradual

change in the grain morphology towards a faceted polygonal morphology was observed,

with equilibrium isotropic tetrakaidodecahedron shape being obtained at 800C (Fig.

5.4(b)). At this temperature the porosity was predominately seen at grain vertexes. At

1000C, pores were seen on the facets of grain boundaries and at grain intersections (Fig.

5.4(c)).

Figure 5.4: Fractured surfaces of ZnO sintered at (a) 600C (b) 800C (c) 1000C.

As the sintering temperature increased, the density increased from 80.5% at 600C to

92.1% at 700 C, while the grain size increased from 166 nm to 443 nm with the increase

in sintering temperature. The relative density increased to 99.4% at 800C with a grain

size of 1.52 µm. Additional rapid grain growth occurred in the 900-1000 C range, with

the grain size increasing to 4.3 µm. This grain growth occurred after maximum relative

density had been achieved (Fig. 5.3b) shows that as the sintering temperature increased

between 600-800C, there was a rapid increase in density with minimal grain growth.

Further increase in sintering temperature resulted in a minor change in the density.

The change in pore morphology may be related to the relative mobility between

grain boundaries and pore as a function of temperature. Typically, at relatively lower

temperatures isolated pores at grain boundaries and grain intersections restrict grain

growth. As the sintering temperature increases, grain boundary mobility exceeds pore

mobility, and the pores are absorbed into the grains. With the grain boundary pinning

no longer in operation, rapid anisotropic grain growth occurs [53]. Formation of intra-

34

granular surface porosity could be a result of closed pores rising to the surface due to

curvature differences between the leading and trailing ends of the pore [54].

5.2.2 Observation of sintering necks

Several locations of the as-sintered disc often contained regions with a liquid-like necks

(Fig. 5.5) suggesting the possibility of a local melting phenomenon. Similar liquid-

like morphology has been previously observed at low homologous temperatures during

SPS processing of Cu [55], Inconel 718 [56] and Mn-Zn ferrites [57]. In the case

of Copper, the necking and liquid-like morphology was attributed to a spark discharge

effect. For Mn-Zn ferrites the liquid like film was attributed to electro-magnetic effects,

while internal current was cited as the reason for the melting of Inconel 718. However,

recent rigorous SPS experiments (using Al2O3 and Cu) remain inconclusive in confirming

the absence/presence of a discharge effect [11] , and the role of electric field and current

has to be further investigated in the case of ionic ceramics.

Figure 5.5: Presence of bridging necks between the particle 600C in (a), (b) and at

1000C in (c).

Further, elemental analysis was carried on the raw powders, sintered samples and in

the neck regions in order to confirm the formation of a new phase or the presence of an

impurity. Elemental maps on the raw powders showed the presence of only Zn and O in

the raw powders and sintered specimens. In the grain boundary regions of the sintered

samples, line scan profiles were taken to look for any compositional gradients. In Fig. 5.6

along the clean boundaries, a uniform composition of Zn and O atoms was observed.

Similar line scans across the sintered necks showed that the composition was ZnO as

seen in Fig. 5.7, further driving the case of high local temperatures.

35

Figure 5.6: Line scan across the grain interfaces showing a dip in Oxygen concentration

in (a) while the composition is uniform in (b) for the sample sintered at 1000C for 0

min.

Figure 5.7: Line scan across the grain interfaces with the necks showing uniform Oxygen

concentration in (a) and (b) for the sample sintered at 1000C for 0 min.

36

Hence the further part of the section shall deal with the theoretical arguments for the

formation of such sintering necks. Thermal gradients can be expected during SPS, both

at the macro scale and at particle interfaces. On the macroscopic scale, the gradients

can arise due to thermal conductivity of the material, radiation effects and the sample

geometry. Finite element models (FEM) have confirmed the possibility of macroscopic

temperature gradients between the edge and the center of the disc in SPS [58]. Melting

of ZnO particles, however, would require temperatures of 2000C, which is unlikely to

occur on the macroscopic scale. Meanwhile, local gradients at the grain boundaries and

particle interfaces can also occur due to preferential Joule heating of the grain boundaries,

dielectric breakdown or the development of a space charge layer [51]. The possibilities

of such field effects in ZnO leading to the observed necks are examined further.

In the present work, the applied field strength was 6.5V/cm−1, whereas techniques like

flash sintering use much larger field strengths. The FEM model developed by Holland

et al. [59]however showed that the local field strength at the particle interfaces can

magnify several times in the case of dielectric materials (due to charge polarization). For

ZnO (dielectric constant =10), a magnification of 30x has been predicted. Narayan [60]

proposed a new mechanism for rapid sintering of ceramics, where accumulation of point

defects along preferred grain boundaries and dislocations increases conductivity. At high

electric field strengths localized melting is possible. Similar localized melting can also

occur in varistor ZnO during electric runaway. Also ZnO treated with pulsed DC current

may cause breakdown by puncture mode and cracking mode, depending on the nature of

the current. In the puncture mode, formation of a through hole from the anode to the

cathode takes place due to vaporization of the ceramic.

Raj et al. [52] have postulated nucleation of Frenkel pair under the influence an exter-

nal field as a fundamental densification mechanism in Field Assisted Sintering Techniques

(FAST). Frenkel pair generation and migration in non-stoichiometric ZnO under electrical

and thermal fields has been reported in the grain boundary defect model for instability in

ZnO varistors [61], where the presence of Zinc interstitials is related to dielectric break-

down. Under thermal and electrical driving forces, intrinsic Zinc interstitials (Znix) are

formed by reaction of defects at the grain boundaries. Accumulation of the neutral Zinc

keeps taking place as long as the electric field is applied, thus increasing conductivity and

ultimately leading to breakdown.

37

From the current observation of necks with a liquid like morphology between sintered

particles (Fig. 5.5), it can be concluded that this phase could in fact aid the densification

and grain growth process. The presence of the liquid layer can enhance the diffusivities

by 6-8 orders of magnitude when compared to solid state sintering, leading to the rapid

densification often expected in SPS experiments. It can be postulated that defect driven

localized currents lead to the observed selective melting of the particle surfaces. Non-

uniform local field strength arising from grain morphology, particle size distribution, stage

of sintering and contamination could be some of the factors leading to the evolution of

the liquid film.

5.2.3 Grain Coalescence

Presence of liquid film possibly enabled densification by nano grain rotation in the samples

sintered at 600C, 800C and 1000C (Fig. 5.5). At 600C, the clusters appeared to be

particle agglomerations (20-30 nm), while at 800C and 1000C clusters were a part of a

well-defined grain (Fig. 5.8). Similar observations of nano grain clusters were observed

during SPS of Al2O3 [62] and Yttrium Aluminum Garnet (YAG) [63].

Figure 5.8: Growth of nano grain clusters from (a) 600C to (b) 800C and (c) 1000C.

Cluster formation behavior can be explained by the grain rotation mechanism in nano

crystalline ceramics [29]. Dense nano clusters (Fig. 5.8(b,c)) can form via a coalescence

route when individual nano grains with high angle grain boundaries rotate until they form

low angle grain boundaries with adjacent nano grains. This mechanism is in contrast

to the traditional grain growth by grain boundary migration, where larger grains grow

at the expense of smaller grains. The relative probability for grain rotation is inversely

38

proportional to the fourth degree of the grain size. Also the extent of subsequent rotation

decreases with each rotation step. Hence, only a fraction of the starting powder particles

are expected to undergo such a densification mechanism, while the rest of the larger

particles undergo curvature driven grain boundary migration [64]. At all temperatures,

there is competition between grain boundary rotation and grain boundary migration.

As the sintering temperature increased the nano grains grow and the tendency for grain

boundary migration increases while that of grain rotation decreases, ultimately leading

to the formation of a single larger grain.

Densification assisted by the grain rotation mechanism is believed to be accelerated

by the presence of a liquid layer. The observation of the liquid like-morphology and

the presence of nano-grains, suggests the possibility of liquid phase driven grain rotation

during SPS of ZnO.

5.2.4 Effect of Dwell Time: Sintering Analysis

For the samples sintered at 600C, increase in holding time from 0 to 15 min at 600C

caused the relative density to increase from 78.2% to 84.3%, while the change in grain

size was statistically insignificant (Fig. 5.9). The microstructure at the end of 15 min

also showed a reduced porostiy, possibly due to better particle rearrangement and break

down of any agglomerates.

Figure 5.9: (a) Change in relative density and (b) grain size as a function of dwell time

at 600C, 700C, 800C and 1000C.

A stress exponent l of 1.24 was calculated from the slope of the plot (Fig. 5.10(a)),

suggesting that the densification was aided by diffusional process, either lattice diffusion,

39

grain boundary diffusion or by viscous flow. From the presence of the liquid morphology

and the formation of nano grain clusters, it is expected that the ZnO particles which are

coated with the viscous layer underwent densification via diffusion through this medium.

Sintering at this temperature was seen to be driven by the formation of nanograin clusters,

the development of necks and minimal grain growth.

The plot of lnG vs lnt (Fig. 5.10(b)) enabled the calculation of the grain growth

exponent for densification. It is generally accepted that the value of the grain growth

exponent characterizes the nature of the grain growth behavior during sintering. At

700C, n was 2.13, while at 1000C and 800C, n was 2.63 and 2.71 respectively. These

values of suggest that the grain growth primarily occurred via volume diffusion through

a liquid layer (n= 3). Previous investigators have also reported a value of n=3 in the

case of undoped ZnO [65]. Such a value can be justified by the formation of the liquid

layer at the particle interfaces, across which migration of Zn2+ ions takes place.

Figure 5.10: (a) Calculation of stress exponent and (b) grain growth exponents.

At 700C sintering temperature, as the holding time increased to 15 min, the relative

density increased from 85.4% to 97%. At 0 min, rounded grains were present. At the end

of the 5 min sintering cycle, the porosity had changed to closed and well defined grains

started to appear (Fig. 5.11(a)). During this time there was a rapid increase in density,

but a slow increase in grain size (Fig. 5.9). Pores at grain quadrature also decreased in

size with increase in holding time. For a sample sintered for 15 min at 700C only a few

pores remained, as seen in Fig. 5.11(a).

At 800C sintering temperature, evolution of pores and grain morphology reached a

steady state. At a dwell time of 10 min, a nearly pore free microstructure was obtained

(Fig. 5.11(b)). At 1000 C fractured surfaces of samples revealed intra-granular surface

40

porosity and pores at grain intersections (Fig. 5.4(c)). The size of these pores increased

with dwell time (Fig. 5.12).

Figure 5.11: (a) SEM images of fractured surfaces of samples sintered at (a) 700C with

15 min dwell and (b) 800C for a dwell time of 10 min.

Figure 5.12: Pore enlargement with increase in dwell time at 1000C.

An increasing negative shrinkage was observed from the ram displacement data for

all the samples sintered at 1000C (Fig. 5.3(b)). A decrease in density after achieving

maximum density has been previously linked to the entrapment of inert gases in the

pores, where the increased gas pressure with an increase in temperature leads to pore

enlargement. However, no appreciable decrease in relative density was observed for any

of the samples sintered at 1000C (Fig. 5.9(b)). An additional sample sintered at 1200C

exhibited further negative shrinkage (from ram displacement data), while its density

41

remained approximately constant. The cause for the observed trend however needs further

investigation. Recently, Olevsky suggested that local temperature gradients due to pulsed

frequencies during SPS can cause thermal diffusion, which drives vacancies and atom

separation. He speculated that the early stages of this mechanism, necking would be

enhanced, but at the final stages, pores act as vacancy sinks under a diffusion gradient,

leading to pore coarsening and hindrance of the densification process. In the present

work, 700 and 800C, shrinkage of pores due to vacancy annihilation could have taken

place while the predicted trend reversal possibly occurred at 1000C.

5.3 Conclusions

Sintering behavior of ZnO prepared through the SPS route was characterized by studying

the microstructural changes and correlating them with possible densification mechanisms.

An increase in relative density was observed up to 86%, without significant grain growth.

Thereafter densification was associated with slow grain growth thereafter. Closure of

inter-granular grains took place, resulting in the achievement of near pore free microstruc-

ture at 800C. Further increase in sintering temperature led to increase in intragranular

surface pore size and negative shrinkage behavior.

Holding time was found to be more significant at lower SPS temperatures in aiding

densification. A window between 600-700C was identified, where the time was most

influential. In this temperature range interplay of sintering parameters controlled the

final porosity and grain size.

The evidence from microstructural analysis supported by literature suggests the pos-

sibility of local melting. Formation of necks with the liquid morphology indicated that

the densification and grain growth was assisted by diffusion through this medium. This

film may also have enabled dense nano grain clusters which contributed to densification

by grain rotation. These observations were supported by sintering analysis exponents (l,

n) which suggested a liquid phase aided densification process.

The significance of electric fields cannot be over emphasized as the effect of pressure;

pulsed current other possible mechanism have not been accounted for. However, the

microstructural evidence indicates that a melting event could have taken place.

42

Chapter 6

Results and Discussion: Spark

Plasma Sintering of MgO

6.1 Introduction

MgO is a well-investigated oxide because of its use as sintering aid and as a refractory.

Recently, transparent MgO ceramics were produced using hot pressing with LiF as an ad-

ditive [66]. Such transparent material are potential replacements for the sapphire windows

used in shock wave experiments. By sintering in vacuum, Misawa et al. have produced

transparent MgO at 1600C. For such applications, a near pore free microstructure is

required and SPS has become a proven technique for producing such pore free ceramics.

Chaim and co-workers have developed transparent MgO using SPS previously and also

have developed detailed densification maps for the sintering at high pressures [9]. Using

the grain growth kinetics, various mechanism were deducted in 900-1420C range [67].

But in light of the recent work by J. Narayan [19] [18], it becomes crucial to investigate

the role of each of the parameters during SPS of nano Mgo.

6.2 Results and Discussion

6.2.1 Decomposition of Mg(OH)2

The as-received powder had a significant amount of Mg(OH)2, as indexed in the XRD

pattern. While it is possible to carry out SPS experiments with this powder by allowing

43

the in-situ decomposition during SPS, the volume change during the phase transfor-

mation is significant to cause sudden displacements and large changes in the estimated

dimensions. Hence the powders were loaded into a muffle furnace and the annealing time

was optimized as 375C for 1h, as the decomposition reaction to form MgO occurs at

332C,as shown in Fig. 6.1.

Figure 6.1: XRD patterns of the heat treated raw powders at various temperatures

After the heat treatment, the powders were deposited on a carbon grid and loaded

into a TEM. The powders were agglomerated with an irregular morphology as shown in

Fig. 6.2. An average crystallite size of 12 nm was calculated. All further SPS experiments

were conducted on this heat treated powder.

Figure 6.2: Bright field image of the heat treated powder with the inset showing the ring

pattern.

44

6.2.2 Grain growth and Densification

The shrinkage curves obtained from the SPS machine for all the samples have been

compared at 1200C. The shrinkage curves for the samples sintered at 50 MPa, showed 3

distinct trends after the initial pressure application at the end of 4 minutes. The shrinkage

increased linearly with temperature till until a slight change of slope was observed in 1050-

1100C temperature range. This was followed by a negative shrinkage as marked in the

Fig. 6.3. However the samples did not show a significant decrease in the density value

that would have corresponded to the observed shrinkage as indicated. A change in the

slope of the shrinkage curve indicates a change in the sintering mechanism.

Figure 6.3: Shrinkage curves of the samples sintered at 1200C with and without the

application of pressure.

The shrinkage curves for both the pressureless samples, with and without the BN

casing showed a similar trend for all the samples. Here also, after the initial shrink-

age corresponding to particle rearrangement and the breakage of soft agglomerates, the

shrinkage increased linearly till the 1150C window discussed above. Here, a prominent

second sintering step was observed that was not observed clearly in the case when pressure

was applied.

The grain sizes estimated from the SEM images and the relative density (theoretical

density of MgO as 3.58 gcm−3) is plotted in Fig. 6.4. The areas marked show that within

experimental error the grain sizes and relative density of the sintered samples in that

window, whether BN layer was used or otherwise, follow a similar trend.

45

Figure 6.4: Comparative (a) grain size and (b) relative density plot of sintered MgO with

the clouds showing similar trends.

At 1100C, for the 50 MPa case, the grain size in the conventional case was 3.01

µm, while in the field blocking experiment it was 3.77 µm. The corresponding relative

densities were 96.32% and 97.51%. For the pressure case, the grain size in the both the

cases was comparatively lower, with 1.04 µm in the conventional case and 1.13 µm in the

BN case. The relative densities of both these samples were similar, with 89.1% for the

BN case and 88.1% for the other.

At 1200C, in the pressure assisted sintering case, the average grain size was 7.14 µm

while for the BN case it was 7.32 µm. The relative densities were 96.92% and 96.53%

respectively. In the pressureless case, for the BN case, the grain size was 3.579 µm while

in the other case it was 3.017µm . The relative density in the case of the BN sample was

92.81% while for the conventional sample it was 91.67%.

As marked in the Fig. 6.4, there are distinct regions in each of the plots. Pressure

application has improved the density of the sample to the window of 95-98% from the 89-

92% window in the pressure-less case. This is an expected behavior as the external driving

force increases, the sinterabilty of the material increases. Also the samples sintered at

higher loads always had a higher grain size. This may be misleading as the relative

densities are not comparable, and 92$ is the commonly suggested density value for the

initiation of the third stage of sintering that involves rapid grain growth. More crucially,

the samples sintered at a given pressure, whether a BN layer is used or not fall in the

same window as seen in Fig. 6.4. This implies that the role of BN seems to be limited in

the set of sintering parameters investigated in this study.

46

Further, the rapid heated sample at 200Cmin−1 showed a lower density at 84.4%

when a 50 MPa load was applied. All the rapidly heated samples crumbled after sintering,

although the microstrucuture shown in Fig. 6.5 had no open porosity that would have lead

to any cracking. The cracking of the sample probably occurred due to the in-homogeneous

thermal gradients generated as a result of the rapid heating.

Figure 6.5: Fracture surface of the sample sintered at 1200C with a heating rate of

200Cmin−1 and a pressure of 50 MPa.

6.2.3 Microstructural Characteristics

Microstructural evolution was tracked for all sample by investigating the fracture surfaces.

The microstrucutres of all the sintered samples was characterized by the presence of

sintering necks in the early stages and terrace-like features on the exposed grain surfaces

in the later stages, as seen in Fig. 6.6(a,b). The necking behavior falls in line with the

classical theory of diffusion driven sintering in the initial stages of sintering. Infact, most

of the microstructural evolution follows the changes that were observed in the case of

ZnO (Refer previous chapter), albeit with a shift in the temperature.

For the 5 MPa samples, at 1100C, the grains were still not faceted with the open

porosity closing down, as observed in Fig. 6.7(a,b). With an increase in sintering tem-

perature to 1300C, the tetrakaidecahodron grains develop with the porosity present at

grain boundaries(Fig. 6.7(c,d). When sintered at 1100C with 50 MPa, the particles

had already become faceted in both the regions as seen in Fig.6.8(a,b). For the sample

47

sintered at 1300C with a pressure of 50 MPa, the pores could be observed entrapped

within the matrix indicating that the grain boundary mobility had become more than

the pore mobility and the final stage of sintering is in process. (Fig. 6.8(c,d)).

Figure 6.6: Microstrucutural feature of the sintered pellet at 1300C with BN (5 MPa)

showing the neck formation while (b) was sintered at 1300C (50MPa) showing a stepped

surface.

Figure 6.7: Comparison of the microstructure of the sintered samples at 5 MPa with

(a),(b) sintered at 1100C and (c),(d) at 1300C with and without BN respectively.

48

Figure 6.8: Comparison of the microstructure of the sintered samples at 50 MPa with

(a),(b) sintered at 1100C and (c),(d) at 1300C with and without BN respectively.

In addition to the similarity of the shrinkage curves, the grain sizes and the relative

density, the microstuctural feature are remarkably similar between the samples sintered at

the same pressure. As seen in Fig. 6.7(a,b), the morphology of the grains is still spherical

in both the cases and evolves to the equilibrium morphology at the same temperatures.

All these observations question the role of electric fields during SPS. With no signif-

icant differences in the microstructure, grain size or the relative density, the utility of

SPS maybe be limited to the application of high loads and rapid heating rates. While

the results of the field experiments by Narayan and co- workers demonstrated the role of

electric field during field treatment of MgO, the fields applied are significantly large than

those commonly used in SPS. By testing the role of the electric field at high pressure,

where plastic deformation is expected to occur and at low temperature, the results indi-

cate that the electric field does not couple with pressure to enhance the sintering rate in

the investigated temperature and pressure regimes.

49

6.2.4 Terraced Oxide Growth

While the presence of steps has been previously attributed to plastic deformation, in the

temperature regime 773-1173K, the yield strength of MgO follows the equation given

below [9],

σy = −0.215T + 314.26

At 1173K the yield strength calculated is 64 MPa. Hence for the temperature range of

1373-1673K investigated in this study it is assumed that the yield strength shall drop

below 50 MPa. Hence, if the presence of terrace features (Fig. 6.9)are attributed to the

plastic deformation, the presence of such structures in the pressure-less experiments at 5

MPa does not fit into this picture.

Figure 6.9: Fractured surface of the sample sintered at 1100C with 5MPa showing

terraced grain structure and (b) showing a pyramid-like terraced structure.

The pyramidal-like shaped terraces observed on the grain surface can also be due to a

vapor-solid process. Using Atomic Force Microscopy Maestre et al. [68] investigated the

growth of such planes in SnO2 samples. The growth such pyramidal features depends on

the crystallographic orientation, surface defects.

Rod Growth During SPS

As seen in Fig. 6.10, large micron sized elongated structures were observed to grow from

the existing grain surfaces and grain boundaries. The nucleation point of such structures

was observed at the tips, at higher magnifications it was observed that the rods also grew

50

from the terrace structures. Maestre et al. [69] had observed similar elongated structures

during sintering of TiO2 along with the observation of the terraced structure. They have

observed that the rods and terraces grew when the sample was subjected to different

sintering cycles. The mechanism for such a growth is a vapor-solid process in which the

source and the substrate for the growth is the compacted grain itself. The nucleation

site for these elongated structures are the single nucleation sites present on the tip of the

terraces.

Figure 6.10: (a) Micron sized rod-like structures while (b)shows the magnified image of

a smaller rod growing from the tip of the terrace structure. The sample was sintered at

1100C with 5 MPa pressure.

Another feature that was observed on the surface of the grains were triangular island

shaped structures. On the fractured surface of the grains, prismatic structures were

decorated on the grain boundaries as seen in Fig. 6.11(a). These structures also have

a terrace-like structure (Fig. 6.11(b)). Whether these are the initial stages of the rod

growth needs to be still investigated. Previously it has been reported that oxidation of a

Cu-5%Pt alloy and the subsequent growth of the Cu2O phase occurred via such an island

nucleation [70]. These triangular protrusions are expected to grow by consuming the local

terrace atoms. Using first principle density functional slabs Jennison and Bogicevic [71]

have shown that even species like (OH)−1 can destabilize the bonds on the surface and

cause nucleation of islands even at elevated temperatures. In the case of SPS of MgO, the

presence of residual Mg(OH)2 can release the hydroxide ions that can aid in the growth

of such structures.

51

Figure 6.11: (a) Triangular islands growing on the surface on the existing micron sized

grains,(b) magnified image of such a structures.The sample was sintered at 1300C with

50 MPa pressure with BN.

6.3 Conclusions

The effect of electric field, pressure and heating rate on the densfication and grain growth

of nano MgO was investigated. It has been observed the pressure helped in lowering the

sintering temperature, while the electric field did not have any significant effect on either

the densification or the grain growth. The trends observed at a pressure of 5 MPa and

50 MPa were same, regardless of the presence of an electric field. Heating of 200Cmin−1

led to a decrease in the density.

Elongated structures were seen to have grown during the sintering process. The pres-

ence of pyramidal terrace structures and triangular structures can aid the nucleation of

such micro and nano rods. The final stage of sintering is characterized by an evaporation-

condensation type mechanism. The role of defect structures, especially oxygen vacancies

needs to be investigated using spectroscopic techniques to fully understand the growth

mechanism.

52

Chapter 7

Conclusions

While the conclusions for each of the materials have been summarized independently,

here a comparative approach is undertaken. Based on the results obtained across the

three oxide sintering experiments, the following conclusions can be drawn:

• In all the three cases, the sintering and grain growth behavior agreed with the

classical sintering theory in the conventional case where pressure was applied. It

was only the use of nanopowders, pressureless sintering experiments and the BN

casing that led to unexpected behavior.

• Pressure assisted sintering always leads to enhanced sintering rates. A lowering of

temperature by atleast 200C was observed to achieve the same density, with all

other parameters remaining the same. Dwell time did not lead to any significant

shrinkage in the case of MgO and NiO. Even in the case of ZnO, only at lower

temperatures, time improved the relative density. Hence shorter dwell times, <3min

can be used to minimize the grain growth.

• The grain growth in nanoparticles is different from micron sized particles as seen

in the case of ZnO and MgO. Coalescence was observed be the preferred pathway

rather than Ostwald Ripening. However, the rapid grain growth to the micron

sized particles needs to be avoided. As seen in the case of NiO, thermally activated

attachment can take place during SPS.

• However, such a behavior was not observed in the case of MgO. MgO and NiO have

the same crystal structure, comparable starting particle size, but the morphologies

observed are different in the case of pressureless sintering.

53

• From the results of MgO and NiO, it appears that the role of electric fields in limited.

In the case of NiO, annealing and SPS had shown similar mechanisms. The BN

experiments also revealed that for MgO, similar trends occur with and without the

fields, while for NiO, a shift in the temperatures was observed in the densification

and grain growth plots. This is likely due to thermal gradients rather than the

electric fields. To conclude this a FEM model with the temperature distributions

needs to be developed. A more quantitative understanding can be developed by

measuring defect driven properties like electrical conductivity.

7.1 Recommendations for Future Work

• A wider range of materials, including metal powders can be investigated for the

role of electric fields. However for a systematic comparative study, all the other

parameters such as particles sizes, plastic deformation need to taken out of the

equation.

• As observed in the case of NiO and MgO, the surface energies of various crystal-

lographic planes seems to play an important role in determining the grain growth

mechanism. Hence, a quantitative thermodynamic description of the surface ener-

gies during sintering needs to be undertaken to predict the growth morphologies.

• Many of the sintering processes lead to the formation of atomistic defects. A quali-

tative description of the defect chemistry using spectroscopic tools will help under-

stand the process better.

• As most of the sintering process occurs in a short period, in-situ microscopy will a

useful tool to monitor the sintering process.

54

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